High temperature alloys

ABSTRACT

Nickel base superalloys in which critical amounts of boron are employed to enhance creep-rupture strength and ductility in the 1,300* F-1,800* F. temperature range. Creep-rupture strength and ductility at temperatures around 1,800* F. also is enhanced by employing amounts of carbon below a critical upper limit. These alloys are particularly useful in the form of castings as gas turbine engine components.

ALLOYING STEELS This application is a continuation of Ser. No. 80,553, filed Oct. 14, 1970, now abandoned.

BACKGROUND OF THE INVENTION 1. Field of the Invention This invention relates to the addition to iron and alloys of iron, hereinafter all referred to as ferrous metals," of an alloying addition, and particularly an alloying addition for the purpose of improving the machinability of the ferrous metals.

2. Description of Prior Art It is well known to add a machinability-improving alloying addition to ferrous metals, but the alloying addition has in general low or negligible solid solubility in ferrous metals so that it is difficult to obtain a good recovery of the alloying addition in the ferrous metal while ensuring that the part of the alloying addition which is retained in the ferrous metal is present in the form of uniformly distributed fine particles.

The presence of large globules of the alloying addition or of a heavy segregation thereof which are commonly found in the bottom of the ingot means that a larger part of the bottom of the ingot must be discarded than the usual discard at the bottom of a rolled ingot.

Furthermore, if the alloying addition is intended to improve the machinability of the ferrous metal, it is important that the alloying addition be dispersed as uniformly as possible throughout the ferrous metal ingot so as to improve the machinability substantially uniformly throughout the ingot as well as avoiding the Iocalised segregation which would be liable to impair the mechanical properties of the ferrous metal.

Also. the fumes generated when the alloying addition is added to ferrous metals may be toxic, in which case the fumes may be efficiently removed throughout the casting process.

The alloying addition may, for example, include one, some or all of the elements lead, selenium and tellurium. in the form of pure metals or in the form of alloys, or as mineral compounds.

SUMMARY OF THE INVENTION An object of the invention is the provision of an improved method of adding the alloying addition to ferrous metals, whereby a high recovery of the alloying addition is obtained in the ingot and an ingot is produced in which the alloying addition is present in the form of a uniformly distributed and finely divided mi cro-dispersion, with a reduction in the segregation of the alloying addition as compared with most commonly used methods so that a reduced proportion of the ingot need be discarded.

A further object of the invention is the provision of an improved method as described above wherein any toxic fumes generated by the alloying addition may be efficiently and easily removed.

According to the present invention we provide a method of adding alloying addition to a ferrous metal, the method comprising the steps of; passing a gas through the ferrous metal, the ferrous metal being heated to a temperature above its customary teeming temperature, the passage ofthe gas being such that turbulence is created in the ferrous metal and a turbulent zone at the surface thereof; simultaneously adding said alloying addition to the ferrous metal at or adjacent to the turbulent zone whilst removing any toxic fumes generated; continuing to pass the gas through the ferrous metal until the temperature of the metal reaches its customary teeming temperature; and'teemin'g the metal into moulds whilst removing any toxic fumes generated by the alloying addition.

The ferrous metal may be tapped from a furnace into a tapping ladle and the gas may be passed through the ferrous metal whilst the metal is in the tapping ladle.

Preferably, the alloying addition is added slowly and in a finely divided form.

Preferably also, the gas is an inert gas such as argon which will not affect the composition of the metal.

The term ferrous metal as used herein includes alloy and other steels and the method of this invention is particularly applicable to ferrous metals of which the carbon content does not exceed 2.0 percent.

Up to 0.50 percent lead may be added to such ferrous metals.

Lead has negligible solid solubility in steel, but is slightly soluble in liquid steel; the solubility of lead increases as the temperature of the steel rises.

Steel is normally tapped from the furnace and cast at temperatures as near its solidifcation temperature as possible since the higher the teeming temperature the more prone is the ingot to defects. It has been found if lead is added to the steel during casting, or in the mould, the temperature of the steel is too low for adequate solubility of the lead, and the ingot tends to contain large globules of lead and to have heavy segregation of lead towards the bottom of the ingot.

Further, the fumes generated when lead is added to steel, are highly toxic and must be efficiently removed throughout the casting process.

Similar problems are encountered with other alloying additions.

Methods hitherto used for adding a machinability improving alloying addition, for example lead, to ferrous metals, for example, plain carbon steel, have obtained recovery of 15 percent to 64 percent lead in the ingot and required approximately 25 percent of the ingot to be discarded because segregation of lead particles and normal metallurgical waste.

The use ofthe method of this invention for the addition of lead to a steel can result in a recovery of about percent lead and requires only the normal ingot discard to be taken to remove pipe and surface defects. Similar advantageous results are obtained with other alloying elements.

BRIEF DESCRIPTION OF THE DRAWINGS One embodiment of the invention will now be described with reference to the accompanying drawings wherein:

FIG. I is a diagrammatic side elevation, partly in section, of a furnace and ladle for use in carrying out the method of the present invention,

FIG. 2 is a diagrammatic cross-sectional view. to an enlarged scale, of the ladle of FIG. 1 showing it in operative position with a fume hood,

FIG. 3 is a diagrammatic cross-sectional view of the apparatus shown in FIG. 2 but with the cross-section taken so as to show the argon inlet and pressure gun for adding lead,

FIG. 4 is a view, to an enlarged scale, of the part of FIG. 3 enclosed in the circle marked 4,

PATENTEUHAR 4 m5 CREEP RUPTURE LIFE- HOURS 3,869.284 SHEET 2 0F 4 EFFECT OF BORON ON THE CREEP RUPTURE STRENGTH OF LOW CARBON ALLOY C NORMAL ALLOY C STRENGTH A. |400F- 94,000 psi B. I800F 29,000 psi I400F- 94,000 psi l800F29,000psi V BORON CONTENT- WT PERCENT PATENTED 43975 3.869284 SHEET 3 0; A

EFFECT OF BORON ON THE CREEP RUPTURE DUCTILITY OF L0w' CARBON ALLOY c.

9- NORMAL ALLOY c DUCTILITY A |400F 94,000psi B I800F 29,000psi 8 '2 I800F 29, 000 psi g 2 2 E3 0 LIJ 4 O Q Hi 95 BORON CONTENT- WT. PERCENT PAIENTEB II/IR 4 I875 SHEET I UP 4 FIG. 5

FIG. 4

FIG?

FIG. 6

HIGH TEMPERATURE ALLOYS FIELD OF THE INVENTION This invention relates to nickel-base alloys having relatively great tensile strength at high temperatures and to castings made from such alloys. The nickel-base superalloys of the present invention are particularly useful for fabricating components of gas turbine engines, such as turbine blades, turbine vanes, integral wheels, and the like.

BACKGROUND OF THE INVENTION There are a number of precipitation strengthened nickel-base superalloys which, because of their strength at high temperatures, are used as materials in fabricating components for use in high temperature sections of gas turbines. The precipitate involved is an intermetallic compound, generally referred to as gamma prime, having the generic formula Ni (Al, Ti). Alloys hardened by such precipitates are referred to as gamma prime strengthed superalloys. In recent years, while characteristics of such alloys at lower temperatures have not altogether been ignored, the greater emphasis in the development of improved alloys has been centered around performance at high'temperatures. High temperature performance has been of concern because of the fact that in new engine designs, as gar turbine operating temperatures are increasing to meet the demands for higher efficiency and power output. High temperature performance properties of particular concern include stress-rupture and creep strength, resistance to thermal fatigue, and corrosion resistance.

It is known that thermal fatigue properties are associated with intermediate temperature (1,3OF l,500F) ductility. The greater the ductility in this range, the more resistant the alloy is to thermal fatigue. As a general rule, alloys with high temperature rupture and creep strength have inadequate thermal fatique and hot corrosion resistance. Conversely, alloys with good hot corrosion resistance show poor high temperature rupture, creep and thermal fatigue properties.

While much work has been done in the development of precipitation strengthened high temperature superalloys, no alloy has been found to be entirely satisfactory with respect to fulfilling the strength, ductility and thermal fatique requirements needed in gas turbine components. Indeed, in recent superalloy developments, strength improvements obtained through composition modification generally have resulted in reduced ductility. In the same manner, alloys designed for improved ductility or toughness or hot corrosion resistance generally possess inadequate strength.

Superalloys suitable for fabricating gas turbine components desirably possess good creep-rupture strength, i.e., resist excessive creep or rupture for long periods of time while under stress at high temperatures. Such alloys also desirably possess good creep-rupture ductility. i.e., deform uniformly and predictably while under stress at high temperature, rather than crack and fracture. Alloys that lack ductility will tolerate little deformation before the onset of crack nucleation, rapid crack propagation, and failure. Use of a material lacking adequate ductility can result in unpredicatable and catastrophic engine component failure. A characteristic peculiar to the gamma prime strengthened superalloys is that they are subject to a sharp decrease in creep-rupture ductility and tensile strength at temperatures between about 1,300F. and l,500F. The decrease in ductility is commonly referred to as the ductility trough," as ductility is higher at temperatures below l,300F. and above l,50ll)F. It generally has been observed that the higher the strength of an alloy, the more pronounced will be the ductility decrease within the ductility trough" temperature range. An example would be MAR-M200 (U.S. Pat. No. 3,164,465 This alloy possesses adequate strength for most advanced gas turbine engine requirements, but lack of l,400F. ductility in the conventionally cast material precludes its usefulness for turbine components.

To circumvent the low ductility problem while retaining usable high temperature strength, the art, in recent years, has turned to a casting process known as directional solidification. This technique, disclosed in U.S. Pat. No. 3,260,505, eliminates graini boundaries that lay in a direction transverse to the direction of applied stress in the component. While directional solidification eliminates a major cause of low longitudinal creep-rupture ductility, it is an expensive procedure and is therefore used only in specialized cases where cost is not a major concern.

It has also been attempted to circumvent ductility trough problems by introducing hafnium to nickel-base superalloys (see, e.g., U.S. Pat. Nos. 3,005,705; 3,677,746; 3,677,746; 3,677,747; and 3,677,748). The addition of very dense and expensive hafnium imposes higher raw material costs and increases the unit weight of the alloys. Increased weight, of course, is a serious disadvantage in alloys intended for aircraft engine components. As is apparent, the lack of the combination of high temperature creep-rupture strength and ductility remainsa major inadequacy inexisting superalloy compositions. These inadequacies are particularly acute since they impair the usefulness of superalloys for many of their intended applications, i.e., formation of gas turbine components.

The alloys of the present invention have improved high temperature strength and corrosion resistance. These alloys are capable of withstandingprolonged op eration at temperatures up to about 2,000F. or higher, and may be'formed into highly advantageous castings.

In accordance with the present invention, alloy compositions have been discovered which possess unique and unusually high creep-rupture strength and ductility in the polycrystalline (non-directionally solidified) form. Specifically, a previously unrecognized criticality has been discovered in the amounts of two alloying elements (boron and carbon) included in chromium, aluminum, and titanium containing nickel base superalloy compositions.

The desirability of adding boron and carbon to high temperature alloys is well documented in the prior art technical and patent literature. The alloy characteristics generally enhanced by the addition of some boron and carbon include ductility, strength, forgeability and in some cases, castability. The present level of technology in the field of superalloy physical metallurgy does not enable precise definition or explanation of the exact mechanism responsible for this property enhancement. Yet one versed in the art of superalloy development recognizes the necessity for the presence of both elements.

While it is known that the role of both carbon and boron in nickel superalloys is complex and dynamic, some generalizations'can be drawn. Carbon appears in the form of complex carbides which prefer grain boundaries as location sites. Detrimental effects on ductility have been noted with certain grain boundary carbide morphologies. This indicates that carbon should be maintained at low levels. On the other hand, it also has been observed that low carbon content re sults in sharply reduced high temperature creep life. It is generally believed, since carbides exert a significant and beneficial effect on rupture strength at high temperature, that carbon should be part of superalloy composition.

Boron is considered an essential ingredient in superalloys. In superalloys, boron in the form ofcomplex borides, is also located at grain boundaries. Grain boundary morphology of superalloys is significant because high temperature creep and rupture failures initiate at and propogate along grain boundaries. Complex borides at grain boundaries reduce the onset of grain boundary tearing under rupture loading.

Typical cast superalloys of the prior art preferably contain carbon in an amount of about 0.10% to about 0.25% by weight. In typical prior art wrought alloys, the carbon content range is between about 0.03% and about 0.15% by weight. For example, in a commercial alloy known as INCO 713, the carbon content is kept as low as 0.05% by weight. Boron content in over fifty prior art alloys studied, preferably is held between 0.007% and 0.03% by weight of the composition. The very small amount of boron used in these commercial alloys demonstrates the potency of the element in affecting properties.

The present invention is based, in part, on the discovery of an unusual and unexpected improvement, in both 1,400F. creep-rupture strength and ductility of gamma prime strengthened nickel-base superalloys, obtained by increasing boron content up to about twenty times the accepted optimum level. Maintenance of the boron content within this critical range of the present invention not only eliminates the problem discussed earlier, relating to the ductility trough present at temperatures between about l,300 F. and l,500 F., but results in a marked increase in creep-rupture strength at those temperatures.

It has also been discovered, in accordance with the present invention, that by reducing the carbon content to a critical upper limit below the amount generally employed in superalloys, it is possible to both effect the improvement in l,400 F. properties and maintain or improve creep-rupture strength and ductility at temperatures around l,800 F. This aspect of the present invention is important with respect to items such as gas turbine components requiring enhanced properties at both l,400 F. and 1,800 F.

Among the alloys of the prior art which will exhibit enhanced properties by following the teachings of the present invention are those disclosed in U.S. Pat. Nos. 3,310,399; 3,164,465; 3,061,426; and 3,619,182. While many of the alloy compositions disclosed in these patents are similar to, and generically overlap with, the alloys of the present invention, none of these patents disclose, nor do corresponding commercial alloy have. the unusual and surprisingly advantageous properties and characteristics of the alloys of the present invention. This is because the prior art fails to recognize the critical carbon and boron content ranges of the alloys of the present invention. All of the commercial alloys derived from the patents referred to above contain substantially less than the minimum boron content used in the alloys of the present invention. Additionally, while at least some of these patents suggest broad boron content ranges which overlap the boron content range of the present invention, there is no recognition that high temperature properties will maximize in a narrow range within these broadly disclosed ranges.

The alloys of the present invention, which have very good stress rupture life at elevated temperatures, contain required minimum amounts of nickel, chromium, aluminum and titanium. The chromium affords primary corrosion resistance while the remaining components are essential to the formation of the gamma prime intermetallic compound, Ni (Al, Ti), which forms the basic superalloy structure of this invention. The Ni Al, Ti) precipitate lends to these alloys their required high temperature strength, and titanium is an important element in providing the strength properties of the present alloys at both room temperature and at elevated temperatures. The presence of significant amounts of Ti strengthener in the present alloys renders them significantly different in character from lower temperature alloys such as those of U.S. Pat. No. 3,005,704, which excludes Ti from its alloys.

SUMMARY OF THE INVENTION In general terms, the present invention pertains to gamma prime phase strengthened superalloys. These alloys are specifically adapted to be employed in cast shapes under conditions of high stress at high temperature. The invention also concerns cast components for use in gas turbine engines made from such alloys.

The alloys of the present invention are predominantly nickel, i.e., at least 35% nickel, and contain in varying amounts, chromium, aluminum, titanium, and boron. One or more ofthe elements carbon, cobalt, zir conium, molybdenum, tantalum, rhenium, columbium, vanadium, and tungsten may also be included in these alloys. In addition, the alloys of the present invention may contain minor amounts of other elements ordinarily included in superalloys by those skilled in the art which will not substantially deleteriously effect the important characteristics of the alloy or which are inadvertently included in such alloys by virtue of impurity levels in commercial grades of alloying ingredients.

It is a principal object of the present invention to include in the aforedescribed alloys amounts of boron within the range of 0.05% to 0.3% by weight to enhance creep-rupture strength and ductility at temperatures around l,400 F. In accordance with preferred embodiments of the present invention, in addition to maintaining the boron content within the range specified, the carbon content of the alloys is maintained below about 0.05% by weight. By additionally maintaining the carbon content below this critical upper limit, it is possible to effect creep-rupture strength and ductility improvement at temperatures around 1,400F. while, at the same time, maintaining or improving creep-rupture strength and ductility at temperatures around l,800 F.

Table l sets forth a broad range and two different narrower ranges, in terms of percent by weight, of elements employed in the alloys of the present invention. It should be understood that the tabulation in Table 1 relates to each element individually and is not intended to solely define composite of broad and narrow ranges.

Nevertheless, composites of the narrower ranges specitied in Table I represent preferred embodiments.

A particularly preferred alloy composition, in percentages by weight, consists essentially of about 8.0% to about 10.25% chromium, about 4.75 to about 5.5% aluminum, about 1.0% to about 2.5% titanium, about 0.05 to about 0.30% (and more preferably about 0.075% to about 0.2%) boron, up to about 0.17% (and more preferably less than 0.05%) carbon, about 8% to about 12% cobalt, about 0.75% to about 1.8% columbium, about 11% to about 16% tungsten, up to 0.20% zirconium, and the balance essentially nickel and minor amounts of impurities and incidental elements which do not detrimentally affect the basic characteristics of the alloy.

TABLE I elongation against boron content, rather than creep rupture life against boron content.

FIG. 4 is a reproduction of a photomicrograph at a magnification of 300, ofa commercial alloy outside the ambit of the present invention.

FIG. 5 is a reproduction of a photomicrograph, at a magnification of 300, of an alloy (comparable to the alloy ofFlG. 4) within the ambit of the present inventlon.

FIG. 6 is a reproduction of a photomicrograph, at a magnification of 7,000, of the same alloy shown in the photomicrograph of FIG. 4.

FIG. 7 is a reproduction of a photomicrograph, at a magnification of 7,000, of the same alloy shown in the photomicrograph of FIG. 5.

ELEMENT BROAD RANGE NARROWER RANGES Another particularly preferred alloy composition, in percentages by weight, consists essentially of about 7.5% to about 8.5% chromium, about 5.75% to 6.25% aluminum, about 0.8% to about 1.2% titanium, about 0.05% to about 0.30% (and more preferably about 0.075% to about 0.2%) boron, up to about 0.13% (and more preferably less than 0.05%) carbon, about 9.5% to about 10.5% cobalt, about 5.75% to about 6.25% molybdenum, about 4.0% to about 4.5% tantalum, 0.05% to 0.10% zirconium, and the balance essentially nickel and minor amounts of impurities and incidental elements which do not detrimentally affect the basic characteristics of the alloy.

Impurities and incidental elements which may be present in the alloys of the present invention include manganese, copper, and silicon in amounts of not more than 0.50%, sulfur and phosphorus in amounts of not more than 0.20%, and iron in amounts of not more than 2.0%. Impurities such as nitrogen, hydrogen, tin, lead, bismuth, calcium, and magnesium should be held to as low a concentration as practical.

BRIEF DESCRIPTION OF THE DRAWINGS FIG. 1 is a graphical plot of percent creep elongation against time for two alloys, one within the ambit of the present invention, and the other outside the ambit of the present invention.

FIG. 2 is a plot of creep-rupture life in hours against the boron content in weight percent of certain nickel base alloys at both 1,400 F. 94,000 psi and 1,800 F. 29,000 psi. The creep-rupture life at l,400 F 94,000 psi and l,800 F 29,000 psi for commercial alloys similar to the alloy of the plot, but outside the ambit ofthe present invention, is also noted on the plot.

FIG. 3 is similar to FIG. 2 but plots percent creep DESCRIPTION OF EXAMPLES AND PREFERRED EMBODIMENTS The alloys of the present invention, containing boron within the critical range of 0.05% to 0.3% by weight. exhibit enhanced creep rupture strength and ductility in the 1,300 F. to l,500 F. temperature range over prior art gamma prime strengthened nickel base superalloys. Thus the alloys of the present invention are capable of withstanding an applied stress of 94,000 psi at 1,400F without rupture for a time in excess of hours. Further, this improvement in strength and ductility properties in the intermediate temperature range (1,300F. 1,500F) is accompanied by a pronounced beneficial effect on high temperature (above 1,700F.)

thermal fatigue properties. Alloys of the present invention, having improved intermediate temperature strength and ductility, demonstrate great advantage in resistance to high temperature thermal fatigue cracking over alloys containing boron in amounts outside the critical range of the present invention.

Designers of gas turbine engines place great importance upon the selection ofcapable and reliable materials. This is particularly true for.rotating components in large aircraft engines where unpredictable engine component failure could endanger the aircraft and its occupants. One of the more critical components in this class of engines is the hot section or turbine blade. Because of the severe conditions of temperature and stress to which these components are subjected, they must be formed of high strength superalloys.

Usual designs involve the mechanical attachment of turbine blades around the periphery of a wheel or disk which rotates at high speed. In operation, hot gases pass over the airfoil portion of the blades, causing the blades and disk to rotate at high speed. The hot gases raise the metal temperatures and the high rotational speed of the disk imposes stress due to centrifugal loading. The attachment, or root portion of the blade is heated only to moderate temperatures due to the cooling effect of the massive disk. The temperature to which the root section of the blade is heated is frequently in the ductility trough temperature range (1,300 F. to 1,500 E). It is an essential mechanical property of an alloy being used for such blades that it be capable of deforming predictably in the root section at temperatures around l,400 F. while withstanding mechanically imposed strain without cracking, i.e., the

alloy must possess reasonable ductility. The alloys of the present invention, containing boron within the critical range of 0.05% to 0.30% by weight, demonstrate great advantage in strength and ductility in the l,400 F. temperature range over prior art alloys intended for use in turbine blades.

The rotating turbine disk, to which the blade root is attached, also requires high resistance to creep and rupture along with ductility and strength to resist fatigue and crack propagation. Accordingly, the alloys described herein provide enhanced properties desirable in disk alloys.

Manufacturers of small gas turbine engines generally employ an integral wheel rather than an assembly ofindividual disks and blades. These integral wheels, consisting of a single component comprising a disk having radially extending blade airfoils at the disk periphery, is usually manufactured by investment casting. Normal modes of operation for small engines subject such components to rapid heating and cooling. This normal mode of operation results in premature cracking at the disk rim between the blade airfoils, because of low cycle thermal and mechanical fatigue. Since the disk rim in many engine designs operates up to about l,400

F., the alloys ofthe present invention enhance the over all performance of integral wheels.

The formulations of several of the more important prior art alloys which are currently commercially used in turbine engines are tabulated in Table 11. The values tabulated represent the amount of each ingredient present in terms of weight percent. The amount of boron and carbon present in each formulation is considered by the prior art to be approximately optimum. With respect to each of the alloys, designated A, B. C. D, E, and F, the U.S. Patent and commercial designation is indicated in the table.

For purposes of comparison, example alloys, compositionally similar to the commercial alloys of Table 11. but containing boron within the critical range of the present invention, were prepared. Analyses ofthese example alloys (designated A-l, B-l, etc.) are presented in Table 111. Standard cast-to-size test bars (0.25 inch in diameter) of the alloys of Table 11 and the example alloys of Table 111 were prepared by melting and casting under vacuum into shell molds. All example alloy specimens were heat treated under a protective atmosphere at l,975 F. for four hours and then air cooled. The example alloys were also subjected to an aging heat treatment at 1,650 F. for ten hours. Each of the commercial alloys of Table 11 was heat treated in accordance with the practice recommended by the alloy developer.

Table IV shows the comparative creep-rupture strength (as measured by timeto rupture) and ductility, (as measured by prior creep) of both commercial alloys A, B, C, and E, and example alloys A-l, B-l, C-l, C-2, C-3 and E-l. All alloys were tested at l,400 F. under a stress of 94,000 psi.

The data in Table IV shows a very significant improvement in both l,400 F. creep-rupture strength and ductility for alloys having a boron content within the critical range of the present invention. At 0.20 weight percent boron, the properties of Example C-3, although decreasing from Example C-2, still show a marked improvement over Alloy C.

The data set forth in Tables 11 IV demonstrates that the utility of nickel-base superalloys for use in gas turbine engine components in which maximum service temperature does not exceed about l,400 F. is greatly enhanced by increasing boron content to an effective level previously considered excessive.

TABLE 11 A* 13* C D* E* F* C 0.10 0.10 0.15 0.15 0.18 0.21 Cr 8.0 10.0 9.0 9.0 10.0 12.5 C0 10.0 10.0 10.0 10.0 15.0 9.0 W 12.5 10.0 3.9 M0 6.0 3.0 2.5 3.0 2.0 Ta 4.25 7.0 1.5 3.9 Ti 1.0 1.0 2.0 1.5 4 7 4.2 A1 6.0 6.0 5.0 5.5 5.5 3.2 B 0.015 0.015 0.015 0.015 0.015 0.02 Zr 0.10 0.10 0.05 0.05 0.06 0.10

V 1.0 N1 (11 (1) (1) (1) (l) (1) A* 3,310,399 13-1900 D* 3,164,465 MAR-M246 13* 3,310,399 B-1910 E* 3,061,4261N-l00 C* 3,164,465 MAR-M200 F* 3,619,1821N-792 (1)1!alancc TABLE 111 EXAMPLE NO.

A-l B-l C-1 C-2 C-3 E-l 9 V TABLE ..C9r tin2 EXAMPLE N0.

A-l B-l C-1 C-2 C-3 13-] Mo 6.06 3.05 3.01 Ta 4.40 6.75 Ti 1.08 1.12 1.98 2.08 L98 4.80 Al 5.95 6.30 4.99 4.80 4.99 5.33 B 0.10 0.10 0.10 0.13 0.20 0.10 Zr 0.05 0.14 0.06 0.03 0.06 0.06 Cb 1.23 1.20 1.23 V 0.86 N1 (1) (l) (l) (l) (l) Balance TABLE IV Creep Rupture Properties Boron Content HOOP/94,000 psi (wt. 7rl Rupture Life (hr) Prior Creep(7r) Alloy A 0.016 31.0 1.98 Example No:

A-l 0.10 229.6 6.80 Alloy B 0.015 102.1 3.68 Example No.:

3-1 0.10 297.2 8.95 Alloy C 0.015 46.7 0.51 Example No.1

(-3 0.20 245.5 2.35 Alloy E 0.012 26.6 096 Example No.2

El 0.l0 345.0 5.25

Prior creep indicates the last creep reading prior to specimen failure The need for improved higher temperature (greater than 1.700F) creep capability in gas turbine alloys is of comparable importance to effecting an improvement in 1,400 F. creep-rupture strength and ductility. Therefore. the effect of the high boron range upon the creep-rupture properties in the l,700 F. to l,900F. temperature range was studied by conducting creeprupture tests on heat treated standard cast-to-size test bars at 1.800 F. under a stress of 29,000 psi.

Results of that testing show that the high boron levels, demonstrated as being unusually effective for l,400 F. properties, were deleterious to l,800 F. rupture strength. The effect was a weakening of the resistance of all alloys in Table II to creep deformation and a noticeable increase in ductility, i.e., a weaker but more ductile material. For gas turbine components requiring both l,400 F. and l,800 F. creep-rupture strength and ductility. use of the alloys shown in Table ll would involve the unacceptable tradeoff of improved l.400 F. ductility at the expense of decreased l,800 F. strength.

lt has been discovered, in further accord with the present invention, that by reducing the carbon content to a critical upper limit of no more than about 0.05 weight percent, it is possible to both effect the improvement in 1,400 F. properties and approximately maintain, and in some cases improve, creep-rupture strength and ductility at l.800 F. Alloys of the present invention, containing less than 0.05 weight percent carbon are capable of withstanding an applied stress of 29,000 psi at 1,800 F. without rupture for a time in excess of 40 hours.

The low carbon aspect of the present invention is particularly important with respect to turbine components requiring enhanced properties at both 1,400 F. and 1,800 F. As previously noted, properties at around l,400 F. are particularly important with respect to the root sections of turbine blades. However, the hot gases passing across the airfoil portion of the blade raise metal temperatures into the l,700 F. to 1.900 F. temperature range. Accordingly, turbine blades desirably require an alloy having good high temperature properties throughout the temperature range of from about 1,300 F. to about 1.900 F. or higher.

To demonstrate the utility and advantages of the low carbon feature of the present invention. thirty pound heats of example alloys A-2, B-2, (-4 through 13, D-l. E-2 through 9, and F- 1 were prepared by melting under vacuum. Standard test bars (0.25 inch diameter) were cast under vacuum into shell molds and all specimens were heat treated under a protective atmosphere at 1,975 F. for four hours. After air cooling, all specimens were subjected to an aging heat treatment of l,650 F. for ten hours. Analyses for series A,B,D, and F example alloys are shown in Table V. Analyses for series C and E example alloys are shown, respectively, in Tables 1V and VII. In all six series compositions, carbon has been reduced to as low a level as possible using normal master alloys and metals in the preparation of each heat. Such a technique is representative of typical commercial practice. Intentional carbon was added however. where appropriate, to determinate the critical upper limit.

Creep-rupture tests were conducted at 1.800 F. under a stress of 29,000 psi and at l.400 F. under a stress of 94,000 psi on all low carbon example alloys. For comparative purposes, the same tests were conducted on the commercial alloys A,B.C,D,E. and F of Table [1. The commercial alloy test bars were heat treated in accordance with the procedures recommended by the producers to achieve maximum mechanical properties. Creep-rupture data for commerical alloys D and F under these conditions were obtained from technical literature provided by the respec- .tive alloy producers.

The data of Table VII demonstrates the applicability of the present invention to a wide range of superalloys. The four example alloys corresponding to the four commercial alloys designated A,B,D, and F had boron and carbon levels approaching the target compositions, i.e., 0.01 weight percent carbon and 0.10 to 0.12 weight percent boron. The comparative test results between the commercial alloys A, B, D, and F and corresponding series A, B, D and F example alloys set forth in Table VIII shows in all cases that very significant improvements are effected at both 1,400 F. and 1,800 F in rupture life and ductility. The

in improving 1,400 F. properties regardless of carbon level. The l,800 F. results show creep-rupture life increasing with increasing boron to about 0.15 weight TABLE V percent. Above 0.15 weight percent boron, strength 15 falls off slightly. Example alloy C-4 shows very good EXAMPLE rupture life at 1,400 E, but the low level of both boron A4 [)4 F4 and carbon causes low ductility in the l,800 F. test. In addition, the combination oflow boron and low carbon 8 822 9-22 contents causes poor castability and a tendency for c6 13,15 11,76 10. 943 20 castings to crack on cooldown during solidification. w 9.66 4.18 M0 m9 310 2.43 204 The minimum boron required to clrcumvent these Ta 3. 53/0 "50 423 problems in the low carbon alloys 15 about 0.05 welght Ti 0.96 0.99 1.38 3.69 percent.

7 Q 888 8:53 88.3 Comparative results of testing between alloy E and Zr 0.073 0.084 0.062 0.083 25 the respective Series E example alloys is shown in (I) (I) (l) Table X. In these data it is seen that although the 1.40

o 111 8.11.1666 0 F. strength is below the high carbon counterparts TABLE VI EXAMPLE NO.

C-4 C-5 C-6 C-7 C-8 C-9 C-IO C-ll C-l2 C-l3 C 0.011 0.010 0.014 0.012 0.011 0.018 0.018 0.045 0.023 0.033 Cr 9.33 8.33 8.89 8.61 8.64 8.97 8.96 9.50 9.54 10.00 CD 10.66 10.70 10.64 10.66 10.71 10.78 10.60 10.50 10.69 10.54 w 12.41 12.40 12.74 12.84 12.48 12.55 12.41 12.5 11.84 13.15 Ti 1.76 1.78 1.77 1.78 1.75 1.77 1.76 2.0 1.75 1.75 Al 5.65 5.53 5.63 5.80 5.41 5.13 5.15 4.98 4.76 4.76 B 0.02 0.03 0.08 0.14 0.15 0.20 0.235 0.10 0.28 0.39 Zr 0.077 0.075 0.079 0.068 0.074 0.065 0.054 0.060 0.053 0.038 Cb 0.95 0.95 0.92 0.92 0.92 0.91 0.85 1 09 0.88 0.79 Ni (1) (l) (l) (I) (I) (l) (l) (l) (I) (I) (I) Balance TABLE VII EXAMPLE N0.

E-2 E-3 E-4 E5 E6 E7 E8 159 C 0.010 0.008 0.008 0.008 0.010 0.011 0.012 0.012 Cr 8.56 9.10 8.95 9.67 9.87 9.87 10.22 10.05 CO 16.60 16.67 16.62 16.62 16.62 16.86 16.80 16.69 MO 3.01 2.94 3.17 3.06 3.03 3.25 3.33 3.32 Ti 4.89 4.90 4.88 4.74 4.91 4.64 4.64 4.56 Al 5.58 5.71 5.60 5.61 5.63 5.22 5.23 5.23 B 0.018 0.044 0.088 0.090 0.125 0.170 0.180 0.220 Zr 0.079 0.067 0.074 0.071 0.060 0.067 0.069 0.064 v 1.06 1.07 1.07 1.08 1.06 0.996 1.01 1.01 i (1) (I) (I) (1) Balance TABLE VIII Creep-Rupture Properties 1400F/94.000 si 1800F/29,0()0psi Boron Carbon Life Prior Life Final (wt (wt (hr) Creep ('71) (hr) Elong.(/r)

- Alloy A 0.016 0.12 31.0 1.98 53.2 6.0

Example No.:

A-2 0.081 0.014 146.5 7.3 44.8 9.9 A116 B 0.010 0.11 102.1 3.68 50.3 9.3 Example No.:

Creep-Rupture Properties 1400F/94,000psi 1800F/29,000psi Boron Carbon Life Prior Life Final (wt 7r) (wt (hr) Creep (hr) Elong.(7r)

Alloy D 0.015 0.15 120.0 2.2 50.0 5.0 Example No.:

D] 0.084 0.009 432.8 4.3 58.1 4.8 Alloy F 0.02 0.21 62.0 3.5 30.0 11.0 Example No.:

TABLE IX Creep-Rupture Properties 1400F/94,000psi 1800F/29,000psi Boron Carbon Life Prior Life Final (wt (wt (hr) Creep (h Elong.(7r)

Alloy C 0.015 0.15 46.7 0.51 96.8 7.5 Example No.1

G4 0.02 0.011 314.8 3.50 73.6 2.2 OS 0.03 0.010 392.2 2.57 85.2 2.4 C-6 0.08 0.014 448.0 2.36 113.9 4.2 O? 0.14 0.012 452.9 3.53 122.4 6.0 C-8 0.15 0.011 468.6 2.21 128.3 5.8 C-9 0.20 0.018 459.8 2.03 117.1 4.5 C-10 0.235 0.018 458.6 1.57 64.8 4.2 C-11 0.10 0.045 397.2 2.59 92.3 7.0 C-12 0.28 0.023 347.4 2.11 43.0 4.6 C13 0.39 0.033 80.7 1.56 14.7 11.]

TABLE X Creep-Rupture Properties 1400F/94.000psi 1800F/29,000psi Boron Carbon Life Prior Life Final (wt'7r) (W17!) (hr) CrceplVr) 1hr) Elong.(7r)

Alloy E 0.015 0.18 26.6 0.95 41.9 8.5 Example No.1

E-2 0.018 0.010 38.6 2.32 27.0 5.2 E-3 0.044 0.008 68.1 5.44 48.6 11.4 E-4 0.088 0.008 104.2 5.92 41.3 12.3 E-S 0.090 0.008 117.2 5.91 38.1 11.7 13-6 0.125 0.010 174.1 4.86 41.2 13.8 E-7 0.170 0.011 266.3 5.03 36.5 11.9 E-S 0.180 0.012 302.2 4.90 31.9 10.1 13-9 0.220 0.012 357.6 5.50 27.6 11.8

reported in Table IV, the improvement over commercial alloy E is significant. In addition, the l,800 F. properties are maintained within a boron range of about 0.05 to 0.15 weight percent. At 0.22 weight percent boron in Example alloy E9. the l.800 F. strength is about sixty percent that of commercial alloy E.

The creep-rupture data discussed previously and presented in Tables 1V, V, V111, IX and X were developed using standard cast-to-size test bars with a 0.250 inch diameter gage section. To demonstrate that the property enhancement is applicable to turbine components, several turbine blade castings were produced from alloy O7 and specimens cut from those castings. Testing was conducted under the same temperature and stress conditions previously employed and results are present in Table XI. The data show the expected reduction in capability compared to test bar properties, but the level of strength and ductility are exceptionally attractive for specimens machined from turbine component castings.

Another major concern of gas turbine engine build ers in the selection of high temperature materials is the ability of the selected alloy to retain initial or starting properties after long time, high temperature exposure. Example Alloy C-7 cast-to-size test bars were subjected for 1,000 hours and examined microstructurally. No deleterious phase formation was observed and subsequent creep-rupture testing was conducted at 1.400F and 94,000 psi for comparison with the same to creep testing at 1,500F under a stress of 40,000 psi TABLE XII-Continued Creep-Rupture Properties I500F exposure for I000 hours under stress of 40,000psi alloy in the as-heat treated condition. Results shown in Table XII reveal essentially no change in rupture life and an improvement in l,400F ductility.

FIG. 1 shows the creep characteristics of typical Alloy C and one of the example alloy C-7 test bars in the l,400F test. In FIG. 1, percent creep elongation is plotted against time. The improved results obtained with the alloys of the present invention are dramatically demonstrated.

FIGS. 2 and 3 further demonstrate the critical relationship between boron content and strength and ductility. FIG. 2 is a plot of creep rupture life in hours against the boron content in weight percent ofC series, low carbon (less than 0.05% by weight), alloys at both l,400F 94,000 psi and l',800F 29,000 psi. The creep rupture life for commercial alloy C at 1,400F 94,000 psi and 1,800F 29,000 psi for commercial alloy C is noted on the plot at, respectively, points A and B. As is apparent, substantial improvements in creep rupture life are obtained at l,400F by maintaining the boron content within the critical range of the present invention.

FIG. 3 is a plot of percent creep elongation against boron content for C series, low carbon alloys at both l,400F 94,000 psi and 1,800F 29,000 psi. The percent creep elongation for commercial alloy C at both l,400F 94,000 psi and l,800F 29,000 psi is also noted on this plot, respectively, at points A and B. Again substantial improvements are apparent at l,400F. with respect to alloys containing boron within the critical range of the present invention. While the percent creep elongation obtained at 1,800F with alloys within the ambit of the present invention is not as high as that ofthe commercial alloy, highly acceptable levels are achieved.

Metallographic examination was conducted in an attempt to explain the mechanism responsible for the observed property enhancement. FIG. 4 shows the normal microstructure of commercial Alloy C in the as-cast condition at 300 magnifications. The light etching dendrite arms or branch-like areas indicate tungsten segregation. A few titanium rich carbides are visible in the lower center portion of the photomicrograph.

The photomicrograph of FIG 5 also at 300 magnifications, shows the profound microstructural change resulting from the added boron and reduced carbon of example alloy C-7. Reducing carbon to less than 0.02 weight percent frees titanium previously tied up as a stable carbide. The increased available titanium in the alloy results in the formation of gamma-gamma prime eutectic in the grain boundaries, a microstructural effect known to enhance l,400F ductility. The boron addition results in the formation of discrete grain bondary particles, identified by electron-beam micro-probe analysis as an M 8 type boride where M (in the C alloy series) is chromium and tungsten. These grain boundary particles are responsible for restoring l,800F creep-rupture ductility to low carbon alloys.

Electron photomicrographs of commercial alloy C and example alloy C-7, at 7,000 magnifications, are shown, respectively, in FIGS. 6 and 7. FIG. 6 shows, as previously stated to be the general case, borides located at the grain boundaries. In FIG. 7, a boride precipitate within each gamma prime particle may be observed, a phenomenon absent in superalloys of the more conventional compositions. The presence of the very fine boride particles appears to retard dislocation movement through the gamma-prime particles and, in

essence, provides dispersion strengthening for improved resistance to creep deformation at 1,800F. This microstructural effect has not been observed in commercial alloys.

Many of the alloys of the present invention may be extruded and hot forged. Wrought, high strength nickel-base superalloys are generally employed in applications where ductility and fracture toughness in the 1,000F to 1,500F temperature range are'of prime concern. Such applications include gas turbine engine turbine and compressor disks. The series E alloys of the present invention may be hot forged, using conventional techniques, into shaped articles having the characteristics considered to be essential in advanced wrought alloys. For example, alloys E-l and E-5 have responded very satisfactorily to extrusion and forging in the 2,000F to 2,200F temperature range in anticipation of the requirements for advanced wrought disk and blade materials.

The present invention also anticipates the use of powder metallurgy for controlling the size, morphology and distribution of the boride microconstituents previously described.

The invention in its broader aspects is not limited to the specific embodiments shown and described. Departures may be made therefrom within the scope of the accompanying claims without departing from the principles of the invention and without sacrificing its chief advantages.

What is claimed is:

l. A nickel base alloy for use at relatively high temperatures consisting essentially of the following elements in the weight percent ranges set forth:

Elements Percent Chromium 5-22 Aluminum 0.2-8 Titanium 0.5-7 Boron 0.07-0.25 Carbon less than 0.0571 Cobalt 0.00-2 Columbium 0.003 Molybdenum 0.00-8 Tantalum 0.00-l0 Vanadium 0.00-2 Tungsten 0,00-20 Rhenium 0.00- Zirconium 000-] 00 3. The nickel base alloy of claim 1 wherein the boron content is about 0.075% to 0.2% by weight.

4. A cast component for use in a gas turbine engine formed of the alloy of claim 1.

5. The component of claim 4 in which said component is a turbine blade.

6. The component of claim 4 in which said component is a disk.

7. The component of claim 4 in which said component is an integral wheel comprising a disk and turbine blade.

8. A cast component for use in a gas turbine engine formed of the alloy of claim 3.

9. The component of claim 8 wherein said component is a turbine blade.

10. The component of claim 8 wherein said component isa disk.

11. The component of claim 8 wherein said component is an integral wheel comprising a disk and turbine blade.

12. A shaped object of the alloy of claim 1 capable of withstanding an applied stress of 94,000 psi at 1,400 F. without rupture for a time in excess of 120 hours.

13. A shaped object of the alloy of claim 1 capable of withstanding an applied stress of 29,000 psi at 1,800 F. without rupture for a time in excess of 40 hours.

14. The alloy of claim 1 which contains, on a weight basis, about 6.0% to about 17% chromium, about 2% to about 8% aluminum, about 0.75% to about 3% titanium, about 2% to about 17% cobalt, and about 40% to 80% by weight nickel.

15. The nickel base alloy of claim 14 wherein the carbon content is no more than 0.025% by weight.

16. A cast component for use in a gas turbine engine formed of the alloy of claim l4.

17. A cast component for use in a gas turbine engine formed of the alloy of claim 15.

18. The alloy of claim 1 which contains, on a weight basis, about 5% to 12% chromium, about 4% to about 8% aluminum. about 0.75% to about 2.5% titanium, about 5% to about 15.5% cobalt, and about 40% to 80% by weight nickel.

19. The nickel base alloy ofclaim 18 wherein the carbon content is no more than 0.025 by weight.

20. A cast component for use in a gas turbine engine formed of the alloy of claim 18.

21. A cast component for use in a gas turbine engine formed of the alloy of claim 19.

22. A nickel base alloy for use at relatively high temperatures consisting essentially of the following elements in the weight percent ranges set forth:

tics of the alloy, said nickel being present in an amount of from about 40% to by weight.

23. The nickel base alloy of claim 22 wherein the boron content is about 0.07% to about 0.25% by weight.

24. The nickel base alloy of claim 22 wherein the carbon content is no more than 0.025 by weight.

25. A cast component for use in a gas turbine engine formed of the alloy of claim 22.

26. A cast component for use in gas turbine engine formed of the alloy of claim 24.

27. A nickel base alloy for use at relatively high temperatures consisting essentially of the following elements in the weight percent ranges set forth:

the balance of the alloy being essentially nickel and minor amounts of impurities and incidental elements which do not detrimentally affect the basic characteristics of the alloy.

28. The nickel base alloy of claim 27 wherein the boron content is about 0.07% to about 0.25% by weight.

29. The nickel base alloy ofclaim 27 wherein the carbon content is no more than 0.025% by weight.

30. A cast component for use in a gas turbine engine formed of the alloy of claim 27.

31. The component of claim 30 in which said component is a turbine blade.

32. A cast component for use in a gas turbine engine formed of the alloy of claim 29.

33. The component of claim 32 wherein said component is a turbine blade.

34. A shaped object of the alloy of claim 27 capable of withstanding an applied stress of 94,000 psi at 1,400" F. without rupture for a time in excess of hours.

35. A shaped object of the alloy of claim 29 capable of withstanding an applied stress of 29,000 psi at l.800 F. without rupture for a time in excess of 40 hours.

36. A nickel base alloy for use at relatively high tem peratures consisting essentially of the following elements in the weight percent ranges set forth:

Elements Percent 65 the balance of the alloy being essentially nickel and minor amounts of impurities and incidental elements which do not detrimentally affect the basic characteristhe balance of the alloy being essentially nickel and minor amounts of impurities and incidental elements which do not detrimentally affect the basic characteristics of the alloy, said nickel being present in an amount of from about 40% to 80% by weight.

37. The nickel base alloy of claim 36 wherein the boron content is about 0.07% to about 0.25% by weight.

38. The nickel base alloy of claim 36 wherein the carbon content is no more than 0.025% by weight.

39. A cast component for use in a gas turbine engine formed of the alloy of claim 36.

40. A cast component for use in a gas turbine engine formed of the alloy of claim 38.

41. A nickel base alloy for use at relatively high temperatures consisting essentially of the following elements in the weight percent ranges set forth:

the balance of the alloy being essentially nickel and minor amounts of impurities and incidental elements which do not detrimentally affect the basic characteristics of the alloy.

42. The nickel base alloy of claim 41 wherein the boron content is about 0.07% to about 0.25% by weight.

43. The nickel base alloy of claim 41 wherein the carbon content is no more than 0.025% by weight.

44. A cast component for use in a gas turbine engine formed of the alloy of claim 41.

45. The component ofclaim 44 in which said component is a turbine blade.

46. A cast component for use in a gas turbine engine formed of the alloy of claim 43.

47. The component of claim 46 wherein said component is a turbine blade.

48. A shaped object of the alloy of claim 41 capable of withstanding an applied stress of 94,000 psi at 1.400 F. without rupture for a time in excess of hours.

49. A shaped object of the alloy of claim 43 capable of,withsta'nding an applied stress of 29,000 psi at l,800 F. without rupture for a time in excess of 40 hours.

50. A nickel base alloy for use at relatively high temperatures consisting essentially of the following elements in the weight percent ranges set forth:

ELEMENTS PERCENT Chromium Aluminum Titanium Boron Carbon Cobalt C olumbium Molybdenum Tantalum Vanadium Tungsten Rhenium Zirconium OKJIN MI I (T PS ll) .3 5 than 0.05 l

formed of the alloy of claim 50. 

1. A NICKLE BASE ALLOY FOR USE AT RELATIVELY HIGH TEMPERATURES CONSISTING ESSENTIALLY OF THE FOLLOWING ELEMENTS IN THE WEIGHT PERCENT RANGES SET FORTH:
 2. The nickel base alloy of claim 1 wherein the carbon content is no more than 0.025% by weight.
 3. The nickel base alloy of claim 1 wherein the boron content is about 0.075% to 0.2% by weight.
 4. A cast component for use in a gas turbine engine formed of the alloy of claim
 1. 5. The component of claim 4 in which said component is a turbine blade.
 6. The component of claim 4 in which said component is a disk.
 7. The component of claim 4 in which said component is an integral wheel comprising a disk and turbine blade.
 8. A cast component for use in a gas turbine engine formed of the alloy of claim
 3. 9. The component of claim 8 wherein said component is a turbine blade.
 10. The component of claim 8 wherein said component is a disk.
 11. The component of claim 8 wherein said component is an integral wheel comprising a disk and turbine blade.
 12. A shaped object of the alloy of claim 1 capable of withstanding an applied stress of 94,000 psi at 1,400* F. without rupture for a time in excess of 120 hours.
 13. A shaped object of the alloy of claim 1 capable of withstanding an applied stress of 29,000 psi at 1,800* F. without rupture for a time in excess of 40 hours.
 14. The alloy of claim 1 which contains, on a weight basis, about 6.0% to about 17% chromium, about 2% to about 8% aluminum, about 0.75% to about 3% titanium, about 2% to about 17% cobalt, and about 40% to 80% by weight nickel.
 15. The nickel base alloy of claim 14 wherein the carbon content is no more than 0.025% by weight.
 16. A cast component for use in a gas turbine engine formed of the alloy of claim
 14. 17. A cast component for use in a gas turbine engine formed of the alloy of claim
 15. 18. The alloy of claim 1 which contains, on a weight basis, about 5% to 12% chromium, about 4% to about 8% aluminum, about 0.75% to about 2.5% titanium, about 5% to about 15.5% cobalt, and about 40% to 80% by weight nickel.
 19. The nickel base alloy of claim 18 wherein the carbon content is no more than 0.025 by weight.
 20. A cast component for use in a gas turbine engine formed of the alloy of claim
 18. 21. A cast component for use in a gas turbine engine formed of the alloy of claim
 19. 22. A nickel base alloy for use at relatively high temperatures consisting essentially of the following elements in the weight percent ranges set forth:
 23. The nickel base alloy of claim 22 wherein the boron content is about 0.07% to about 0.25% by weight.
 24. The nickel base alloy of claim 22 wherein the carbon content is no more than 0.025 by weight.
 25. A cast component for use in a gas turbine engine formed of the alloy of claim
 22. 26. A cast component for use in gas turbine engine formed of the alloy of claim
 24. 27. A nickel base alloy for use at relatively high temperatures consisting essentially of the following elements in the weight percent ranges set forth:
 28. The nickel base alloy of claim 27 wherein the boron content is about 0.07% to about 0.25% by weight.
 29. The nickel base alloy of claim 27 wherein the carbon content is no more than 0.025% by weight.
 30. A cast component for use in a gas turbine engine formed of the alloy of claim
 27. 31. The component of claim 30 in which said component is a turbine blade.
 32. A cast component for use in a gas turbine engine formed of the alloy of claim
 29. 33. The component of claim 32 wherein said component is a turbine blade.
 34. A shaped object of the alloy of claim 27 capable of withstanding an applied stress of 94,000 psi at 1,400* F. without rupture for a time in excess of 120 hours.
 35. A shaped object of the alloy of claim 29 capable of withstanding an applied stress of 29,000 psi at 1,800* F. without rupture for a time in excess of 40 hours.
 36. A nickel base alloy for use at relatively high temperatures consisting essentially of the following elements in the weight percent ranges set forth:
 37. The nickel base alloy of claim 36 wherein the boron content is about 0.07% to about 0.25% by weight.
 38. The nickel base alloy of claim 36 wherein the carbon content is no more than 0.025% by weight.
 39. A cast component for use in a gas turbine engine formed of the alloy of claim
 36. 40. A cast component for use in a gas turbine engine formed of the alloy of claim
 38. 41. A nickel base alloy for use at relatively high temperatures consisting essentially of the following elements in the weight percent ranges set forth:
 42. The nickel base alloy of claim 41 wherein the boron content is abOut 0.07% to about 0.25% by weight.
 43. The nickel base alloy of claim 41 wherein the carbon content is no more than 0.025% by weight.
 44. A cast component for use in a gas turbine engine formed of the alloy of claim
 41. 45. The component of claim 44 in which said component is a turbine blade.
 46. A cast component for use in a gas turbine engine formed of the alloy of claim
 43. 47. The component of claim 46 wherein said component is a turbine blade.
 48. A shaped object of the alloy of claim 41 capable of withstanding an applied stress of 94,000 psi at 1,400* F. without rupture for a time in excess of 120 hours.
 49. A shaped object of the alloy of claim 43 capable of withstanding an applied stress of 29,000 psi at 1,800* F. without rupture for a time in excess of 40 hours.
 50. A nickel base alloy for use at relatively high temperatures consisting essentially of the following elements in the weight percent ranges set forth:
 51. The nickel base alloy of claim 50 wherein the boron content is about 0.07% to about 0.25% by weight.
 52. The nickel base alloy of claim 51 wherein the carbon content is no more than 0.025% by weight.
 53. A cast component for use in a gas turbine engine formed of the alloy of claim
 50. 